Steel plate to be precipitating tinfor welded structures, method for manufacturing the same and welding fabric using the same

ABSTRACT

A weldable structural steel product having fine complex precipitates of TiN and MnS is provided which contains, in terms of percent by weight, 0.03 to 0.17% C, 0.01 to 0.05% Si, 1.0 to 2.5% Mn, 0.05 to 0.2% Ti, 0.0005 to 0.1% Al, 0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P, 0.003 to 0.05% S, at most 0.005 % O, and balance Fe and incidental impurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, 6.5≦(Ti+2Al+4B)/N≦14, and 220≦Mn/S≦400. The steel has a microstructure consisting essentially of a complex structure of ferrite and pearlite having a grain size of 20 μm or less.

TECHNICAL FIELD

[0001] The present invention relates to a structural steel productsuitable for use in constructions, bridges, ship constructions, marinestructures, steel pipes, line pipes, etc. More particularly, the presentinvention relates to a welding structural steel product which ismanufactured using fine complex precipitates of TiN and MnS dispersed insuch a fashion that MnS surrounds TiN, thereby being capable ofsimultaneously exhibiting improved toughness and strength in aheat-affected zone. The, present invention also relates to a method formanufacturing the welding structural steel product, and a weldedconstruction using the welding structural steel product.

BACKGROUND ART

[0002] Recently, as the height or size of buildings and other structureshas increased, steel products having an increased size have beenincreasingly used. That is, thick steel products have been increasinglyused. In order to weld such thick steel products, it is necessary to usea welding process with a high efficiency. For welding techniques forthick steel products, a heat-input submerged welding process enabling asingle pass welding, and an electro-welding process have been widelyused. The heat-input welding process enabling a single pass welding isalso applied to ship constructions and bridges requiring welding ofsteel plates having a thickness of 25 mm or more. Generally, it ispossible to reduce the number of welding passes at a higher amount ofheat input because the amount of welded metal is increased. Accordingly,there may be an advantage in terms of welding efficiency where theheat-input welding process is applicable. That is, in the case of awelding process using an increased heat input, its application can bewidened. Typically, the heat input used in welding process are in therange of 100 to 200 kJ/cm. In order to weld steel plates furtherthickened to a thickness of 50 mm or more, it is necessary to usesuper-high heat input ranging from 200 kJ/cm to 500 kJ/cm.

[0003] Where high heat input is applied to a steel product, the heataffected zone, in particular, its portion arranged near a fusionboundary, is heated to a temperature approximate to a melting point ofthe steel product by welding heat input. As a result, growth of grainsoccurs at the heat affected zone, so that a coarsened grain structure isformed. Furthermore, when the steel product is subjected to a coolingprocess, fine structures having degraded toughness, such as bainite andmartensite, may be formed. Thus, the heat affected zone may be a siteexhibiting degraded toughness.

[0004] In order to secure a desired stability of such a weldingstructure, it is necessary to suppress the growth of austenite grains atthe heat affected zone, so as to allow the welding structure to maintaina fine structure. Known as means for meeting this requirement aretechniques in which oxides stable at a high temperature or Ti-basedcarbon nitrides are appropriately dispersed in steels in order to delaygrowth of grains at the heat affected zone during a welding process.Such techniques are disclosed in Japanese Patent Laid-open PublicationNo. Hei. 12-226633, Hei. 11-140582, Hei. 10-298708, Hei. 10-298706, Hei.9-194990, Hei. 9-324238, Hei. 8-60292, Sho. 60-245768, Hei. 5-186848,Sho. 58-31065, Sho. 61-797456, and Sho. 64-15320, and Journal ofJapanese Welding Society, Vol. 52, No. 2, pp 49.

[0005] The technique disclosed in Japanese Patent Laid-open PublicationNo. Hei. 11-140582 is a representative one of techniques usingprecipitates of TiN. This technique has proposed structural steelsexhibiting an impact toughness of about 200 J at 0° C. (in the case of amatrix, about 300 J). In accordance with this technique, the ratio ofTi/N is controlled to be 4 to 12, so as to form TiN precipitates havinga grain size of 0.05 μm or less at a density of 5.8×10³/mm² to 8.1×10⁴/mm² while forming TiN precipitates having a grain size of 0.03 to 0.2μm at a density of 3.9×10³/mm² to 6.2×10⁴ /mm², thereby securing adesired toughness at the welding site. In accordance with thistechnique, however, both the matrix and the heat affected zone exhibitsubstantially low toughness where a heat-input welding process isapplied. For example, the matrix and heat affected zone exhibit impacttoughness of 320 J and 220 J at 0° C. Furthermore, since there is aconsiderable toughness difference between the matrix and heat affectedzone, as much as about 100 J, it is difficult to secure a desiredreliability for a steel construction obtained by subjecting thickenedsteel products to a welding process using super-high heat input.Moreover, in order to obtain desired TiN precipitates, the techniqueinvolves a process of heating a slab at a temperature of 1,050° C. ormore, quenching the heated slab, and again heating the quenched slab fora subsequent hot rolling process. Due to such a double heat treatment,an increase in the manufacturing costs occurs.

[0006] Japanese Patent Laid-Qpen Publication No. Hei. 9-194990 disclosesa technique in which the ratio between Al and O in low steel (N≦0.005%)is controlled to be within a range of 0.3 to 1.5 (0.3≦Al/O≦1.5) in orderto form a complex oxide containing Al, Mn, and Si. However, the steelproduct according to this technique exhibits a degraded toughnessbecause when a welding process using a high heat input of about 100kJ/cm, the transition temperature at the heat affected zone correspondsto a level of is about −50. Also, Japanese Patent Laid-open PublicationNo. Hei. 10-298708 discloses a technique in which complex precipitatesof MgO and TiN are utilized. However, the steel product according tothis technique exhibits a degraded toughness in that when a weldingprocess using a high heat input of about 100 kJ/cm, the impact toughnessat 0° C. in the heat affected zone corresponds to 130 J.

[0007] There have been many techniques for improving the toughness ofthe heat affected zone using TiN precipitates and Al-based oxides or MgOwhere a welding process using a high heat input is applied. However,there is no technique capable of remarkably improving the toughness ofthe heat affected zone where a welding process using a super-high heatinput is carried out for a prolonged period of time at 1,350° C. ormore.

DISCLOSURE OF THE INVENTION

[0008] Therefore, it is an object of the invention to provide a weldingstructural steel product in which complex precipitates of TiN and MnSare dispersed in such a fashion that MnS surround TiN precipitates,thereby being capable of improving both the toughness and strength (orhardness) of the heat affected zone while minimizing the toughnessdifference between the matrix and the heat affected zone, a method formanufacturing the welding structural steel product, and a weldedstructure using the welding structural steel product.

[0009] In accordance with one aspect, the present invention provides awelding structural steel product having fine complex precipitates of TiNand MnS, comprising, in terms of percent by weight, 0.03 to 0.17% C,0.01 to 0.5% Si, 1.0 to 2.5% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al,0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P,0.003 to 0.05% S, at most 0.005% O, and balance Fe and incidentalimpurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40,2.5≦Al/N≦7, 6.5≦(Ti+2Al+4B)/N≦14, and 200≦Mn/S≦400, and having amicrostructure essentially consisting of a complex structure of ferriteand pearlite having a grain size of 20 μm or less.

[0010] In accordance with another aspect, the present invention providesa method for manufacturing a welding structural steel product havingfine complex precipitates of TiN and MnS, comprising the steps of:

[0011] preparing a steel slab containing, in terms of percent by weight,0.03 to 0.17% C, 0.01 to 0.5% Si, 1.0 to 2.5% Mn, 0.005 to 0.2% Ti,0.0005 to 0.1% Al, 0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2%W, at most 0.03% P, 0.003 to 0.05% S, at most 0.005% O, and balance Feand incidental impurities while satisfying conditions of 1.2≦Ti/N≦2.5,10≦N/B≦40, 2.5≦Al/N≦7, 6.5≦(Ti+2Al+4B)/N≦14, and 200≦Mn/S≦400;

[0012] heating the steel slab at, a temperature ranging from 1,000° C.to 1,250° C. for 60 to 180 minutes;

[0013] hot rolling the heated steel slab in an austeniterecrystallization range at a thickness reduction rate of 40% or more;and

[0014] cooling the hot-rolled steel slab at a rate of 1° C./min to atemperature corresponding to ±10° C. from a ferrite transformationfinish temperature.

[0015] In accordance with another aspect, the present invention providesa method for manufacturing a welding structural steel product havingfine complex precipitates of TiN and MnS, comprising the steps of:

[0016] preparing a steel slab containing, in terms of percent by weight,0.03 to 0.17% C, 0.01 to 0.5% Si, 1.0 to 2.5% Mn, 0.005 to 0.2% Ti,0.0005 to 0.1% Al, at most 0.005 N, 0.0003 to 0.01% B, 0.001 to 0.2% W,at most 0.03% P, 0.003 to 0.05% S, at most 0.005% O, and balance Fe andincidental impurities while satisfying a condition of 200≦Mn/S≦400;

[0017] heating the steel slab at a temperature ranging from 1,000° C. to1,250° C. for 60 to 180 minutes while nitrogenizing the steel slab tocontrol the N content of the steel slab to be 0.008 to 0.03%, and tosatisfy conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, and6.5≦(Ti+2Al+4B)/N≦14;

[0018] hot rolling the nitrogenized steel slab in an austeniterecrystallization range at a thickness reduction rate of 40% or more;and

[0019] cooling the hot-rolled steel slab at a rate of 1° C./min to atemperature corresponding to ±10° C. from a ferrite transformationfinish temperature.

[0020] In accordance with another aspect, the present invention providesa welded structure having a superior heat affected zone toughness,manufactured using any one of the above described welding structuralsteel products.

[0021] Best Mode for Carrying Out the Invention

[0022] Now, the present invention will be described in detail.

[0023] In the specification, the term “prior austenite” represents anaustenite formed at the heat affected zone in a steel product (matrix)when a welding process using high heat input is applied to the steelproduct. This austenite is distinguished from the austenite formed inthe manufacturing procedure (hot rolling process).

[0024] After carefully observing the growth behavior of the prioraustenite in the heat affected zone in a steel product (matrix) and thephase transformation of the prior austenite exhibited during a coolingprocedure when a welding process using high heat input is applied to thesteel product, the inventors found that the heat affected zone exhibitsa variation in toughness with reference to the critical grain size ofthe prior austenite (about 80 μm), and that the toughness at the heataffected zone is increased at an increased fraction of fine ferrite.

[0025] On the basis of such an observation, the present invention ischaracterized by:

[0026] [1] utilizing complex precipitates of TiN and MnS in the steelproduct;

[0027] [2] reducing the grain size of initial ferrite in the steelproduct (matrix) to a critical level or less so as to control the prioraustenite to have a grain size of about 80 μm or less; and

[0028] [3] reducing the ratio of Ti/N to effectively form BN and AlNprecipitates, thereby increasing the fraction of ferrite at the heataffected zone, while controlling the ferrite to have a acicular orpolygonal structure effective to achieve an improvement in toughness.

[0029] The above features [1], [2], [3] of the present invention will bedescribed in detail.

[0030] [1] Complex Precipitates of TiN and MnS

[0031] Where a high heat-input welding is applied to a structural steelproduct, the heat affected zone near a fusion boundary is heated to ahigh temperature of about 1,400° C. or more. As a result, TiNprecipitated in the matrix is partially dissolved due to the weld heat.Otherwise, an Ostwald ripening phenomenon occurs. That is, precipitateshaving a small grain size are dissolved, so that they are diffused inthe form of precipitates having a larger grain size. In accordance withthe Ostwald ripening phenomenon, a part of the precipitates arecoarsened. Furthermore, the density of TiN precipitates is considerablyreduced, so that the effect of suppressing growth of prior austenitegrains disappears.

[0032] After observing a variation in the characteristics of TiNprecipitates depending on the ratio of Ti/N while taking intoconsideration the fact that the above phenomenon may be caused bydiffusion of Ti atoms occurring when TiN precipitates dispersed in thematrix are dissolved by the welding heat, the inventors discovered thenew fact that under a high nitrogen concentration condition (that is, alow Ti/N ratio), the concentration and diffusion rate of dissolved Tiatoms are reduced, and an improved high-temperature stability of TiNprecipitates is obtained. That is, when the ratio between Ti and N(Ti/N) ranges from 1.2 to 2.5, the amount of dissolved Ti is greatlyreduced, thereby causing TiN precipitates to have an increasedhigh-temperature stability. As a result, fine TiN precipitates areuniformly dispersed at a high density. Such a surprising result wasassumed to be based on the fact that the solubility product representingthe high-temperature stability of TiN precipitates is reduced at areduced content of nitrogen, because when the content of nitrogen isincreased under the condition in which the content of Ti is constant,all dissolved Ti atoms are easily coupled with nitrogen atoms, and theamount of dissolved Ti is reduced under a high nitrogen concentrationcondition.

[0033] Also, the inventors noticed that if re-dissolution of TiNprecipitates distributed in the heat affected zone near the fusionboundary can be prevented even when those TiN precipitates in the matrixare fine while being uniformly dispersed, it is possible to easilysuppress growth of prior austenite grains. That is, the inventorsresearched a scheme for delaying the re-dissolution of TiN precipitatesin a matrix. As a result of this research, the inventors found thatwhere TiN is distributed in the heat affected zone in the form ofcomplex precipitates of TiN and MnS in such a fashion that MnS surroundsTiN precipitates in the matrix, re-dissolution of those TiN precipitatesinto the matrix is considerably delayed even when the TiN precipitatesare heated to a high temperature of 1,350° C. That is, MnS, which ispreferentially re-dissolved, surrounds TiN, so that it influences thedissolution of TiN and the re-dissolution rate of TiN into the matrix.As a result, TiN effectively contributes to suppressing growth of prioraustenite grains. Thus, a remarkable improvement in the toughness of theheat affected zone is achieved.

[0034] Accordingly, it is important to reduce the solubility productrepresenting the high-temperature stability of TiN precipitates whileuniformly dispersing fine complex precipitates of TiN and MnS in thematrix. After observing variations in the size, amount, and density ofcomplex precipitates of TiN and MnS depending on the ratios of Ti and N(Ti/N) and of Mn and S (Mn/S), the inventors found that complexprecipitates of TiN and MnS having a grain size of 0.01 to 0.1 μm areprecipitated at a density of 1.0×10⁷/mm² or more under the condition inwhich the ratio of Ti/N is 1.2 to 2.5, and the ratio of Mn/S is 220 to400. That is, the precipitates had a uniform space of about 0.5 μm.

[0035] The inventors also discovered an interesting fact. That is, evenwhen a high-nitrogen steel is manufactured by producing, from a steelslab, a low-nitrogen steel having a nitrogen content of 0.005% or lessto exhibit a low possibility of generation of slab surface cracks, andthen subjecting the low-nitrogen steel to a nitrogen zing treatment in aslab heating furnace, it is possible to obtain desired TiN precipitatesas defined above, in so far as the ratio of Ti/N is controlled to be 1.2to 2.5. This was analyzed to be based on the fact that when an increasein nitrogen content is made in accordance with a nitrogen zing treatmentunder the condition in which the content of Ti is constant, alldissolved Ti atoms are easily rendered to be coupled with nitrogenatoms, thereby reducing the solubility product of TiN representing thehigh-temperature stability of TiN precipitates.

[0036] In accordance with the present invention, in addition to thecontrol of the ratio of Ti/N, respective ratios of N/B, Al/N, and V/N,the content of N, and the total content of Ti+Al+B+(V) are generallycontrolled to precipitate N in the form of BN, AlN, and VN, taking intoconsideration the fact that promoted aging may occur due to the presenceof dissolved N under a high-nitrogen environment. In accordance with thepresent invention, as described above, the toughness difference betweenthe matrix and the heat affected zone is minimized by not onlycontrolling the density of TiN precipitates depending on the ratio ofTi/N and the solubility product of TiN, but also dispersing TiN in theform of complex precipitates of TiN and MnS in which MnS appropriatelysurrounds TiN precipitates. This scheme is considerably different fromthe conventional precipitate control scheme (Japanese Patent Laid-openPublication No. Hei. 11-140582) in which the amount of TiN precipitatesis increased by simply increasing the content of Ti (Ti/N≧4).

[0037] [2] Control for Ferrite Grain Size of Steels (Matrix)

[0038] After research, the inventors found that in order to controlprior austenite to have a grain size of about 80 μm or less, it isimportant to form fine ferrite grains in a complex structure of ferriteand pearlite, in addition to control of precipitates. Fining of ferritegrains can be achieved by fining austenite grains in accordance with ahot rolling process or controlling growth of ferrite grains occurringduring, a cooling process following the hot rolling process. In thisconnection, it was also found that it is very effective to appropriatelyprecipitate carbides (VC and WC) effective to growth of ferrite grainsat a desired density.

[0039] [3] Microstructure of Heat Affected Zone

[0040] The inventors also found that the toughness of the heat affectedzone is considerably influenced by not only the size of prior austenitegrains, but also the amount and shape of ferrite precipitated at thegrain boundary of the prior austenite when the matrix is heated to atemperature of 1,400° C. In particular, it is preferable to generate atransformation of polygonal ferrite or acicular ferrite in austenitegrains. For this transformation, AlN and BN precipitates are utilized inaccordance with the present invention.

[0041] The present invention will now be described in conjunction withrespective components of a steel product to be manufactured, and amanufacturing method for the steel product.

[0042] [Welding Structural Steel Product]

[0043] First, the composition of the welding structural steel productaccording to the present invention will be described.

[0044] In accordance with the present invention, the content of carbon(C) is limited to a range of 0.03 to 0.17weight % (hereinafter, simplyreferred to as “%”).

[0045] Where the content of carbon (C) is less than 0.03%, it isimpossible to secure a sufficient strength for structural steels. On theother hand, where the C content exceeds 0.17%, transformation ofweak-toughness microstructures such as upper bainite, martensite, anddegenerate pearlite occurs during a cooling process, thereby causing thestructural steel product to exhibit a degraded low-temperature impacttoughness. Also, an increase in the hardness or strength of the weldingsite occurs, thereby causing a degradation in toughness and generationof welding cracks.

[0046] The content of silicon (Si) is limited to a range of 0.01 to0.5%.

[0047] At a silicon content of less than 0.01%, it is impossible toobtain a sufficient deoxidizing effect of molten steel in the steelmanufacturing process. In this case, the steel product also exhibits adegraded corrosion resistance. On the other hand, where the siliconcontent exceeds 0.5%, a saturated deoxidizing effect is exhibited. Also,transformation of island-like martensite is promoted due to an increasein hardenability occurring in a cooling process following a rollingprocess. As a result, a degradation in low-temperature impact toughnessoccurs.

[0048] The content of manganese (Mn) is limited to a range of 1.0 to2.5%.

[0049] Mn has an effective function for improving the deoxidizingeffect, weldability, hot workability, and strength of steels. Thiselement is precipitated in the form of MnS around Ti-based oxides, sothat it promotes generation of acicular and polygonal ferrite effectiveto improve the toughness of the heat affected zone. The Mn element formsa substitutional solid solution in a matrix, thereby solid-solutionstrengthening the matrix to secure desired strength and toughness. Inorder to obtain such effects, it is desirable for Mn to be contained inthe composition in a content of 1.0% or more. However, where the Mncontent exceeds 2.5%, macroscopic segregation and microscopicsegregation occur in accordance with a segregation mechanism in asolidification procedure of steels, thereby promoting formation of acentral segregation band in the matrix in a rolling process. Such acentral segregation band serves as a cause for forming a centrallow-temperature transformed structure in the matrix.

[0050] The content of titanium (Ti) is limited to a range of 0.005 to0.2%.

[0051] Ti is an essential element in the present invention because it iscoupled with N to form fine TiN precipitates stable at a hightemperature. In order to obtain such an effect of precipitating fine TiNgrains, it is desirable to add Ti in an amount of 0.005% or more.However, where the Ti content exceeds 0.2%, coarse TiN precipitates andTi oxides may be formed in molten steel. In this case, it is impossibleto suppress the growth of prior austenite grains in the heat affectedzone.

[0052] The content of aluminum (Al) is limited to a range of 0.0005 to0.1%.

[0053] Al is an element which is not only necessarily used as adeoxidizer, but also serves to form fine AlN precipitates in steels. Alalso reacts with oxygen to form an Al oxide, thereby preventing Ti fromreacting with oxygen. Thus, Al aids Ti to form fine TiN precipitates.For such functions, Al is preferably added in an amount of 0.0005% ormore. However, when the content of Al exceeds 0.1%, dissolved Alremaining after precipitation of AlN promotes formation of Widmanstattenferrite and island-like martensite exhibiting weak toughness in the heataffected zone in a cooling process. As a result, a degradation in thetoughness of the heat affected zone occurs where a high heat inputwelding process is applied.

[0054] The content of nitrogen (N) is limited to a range of 0.008 to0.03%.

[0055] N is an element essentially required to form TiN, AlN, BN, VN,NbN, etc. N serves to suppress, as much as possible, the growth of prioraustenite grains in the heat affected zone when a high heat inputwelding process is carried out, while increasing the amount ofprecipitates such as TiN, AlN, BN, VN, NbN, etc. The lower limit of Ncontent is determined to be 0.008% because N considerably affects thegrain size, space, and density of TiN and AlN precipitates, thefrequency of those precipitates to form complex precipitates withoxides, and the high-temperature stability of those precipitates.However, when the N content exceeds 0.03%, such effects are saturated.In this case, a degradation in toughness occurs due to an increasedamount of dissolved nitrogen in the heat affected zone. Furthermore, thesurplus N may be included in the welding metal in accordance with adilution occurring in the welding process, thereby causing a degradationin the toughness of the welding metal.

[0056] Meanwhile, the slab used in accordance with the present inventionmay be low-nitrogen steels which may be subsequently subjected to anitrogen zing treatment to form high-nitrogen steels. In this case, theslab has a N content of 0.0005% in order to exhibit a low possibility ofgeneration of slab surface cracks. The slab is then subjected to are-heating process involving a nitrogen zing treatment, so as tomanufacture high-nitrogen steels having an N content of 0.008 to 0.03%

[0057] The content of boron (B) is limited to a range of 0.0003 to0.01%.

[0058] B is an element which is very effective to form acicular ferriteexhibiting a superior toughness in grain boundaries while formingpolygonal ferrites in the grain boundaries. B forms BN precipitates,thereby suppressing the growth of prior austenite grains. Also, B formsFe boron carbides in grain boundaries and within grains, therebypromoting transformation into acicular and polygonal ferrites exhibitinga superior toughness. It is impossible to expect such effects when the Bcontent is less than 0.0003%. On the other hand, when the B contentexceeds 0.01%, an increase in hardenability may undesirably occur, sothat there may be possibilities of hardening the heat affected zone, andgenerating low-temperature cracks.

[0059] The content of tungsten (W) is limited to a range of 0.001 to0.2%.

[0060] When tungsten is subjected to a hot rolling process, it isuniformly precipitated in the form of tungsten carbides (WC) in thematrix, thereby effectively suppressing growth of ferrite grains afterferrite transformation. Tungsten also serves to suppress the growth ofprior austenite grains at the initial stage of a heating process for theheat affected zone. Where the tungsten content is less than 0.001%, thetungsten carbides serving to suppress the growth of ferrite grainsduring a cooling process following the hot rolling process are dispersedat an insufficient density. On the other hand, where the tungstencontent exceeds 0.2%, the effect of tungsten is saturated.

[0061] The content of phosphorous (P) is limited to 0.030% or less.

[0062] Since P is an impurity element causing central segregation in arolling process and formation of high-temperature cracks in a weldingprocess, it is desirable to control the content of P to be as low aspossible. In order to achieve an improvement in the toughness of theheat affected zone and a reduction in central segregation, it isdesirable for the P content to be 0.03% or less.

[0063] The content of sulfur (S) is limited to a range of 0.003 to0.005%.

[0064] S is an element which is precipitated around Ti-based oxides inthe form of MnS, so that it influences the formation of ferrites havingan acicular or polygonal structure effective to achieve an improvementin the toughness of the heat affected zone. For such effects, S ispreferably added in an amount of 0.003% or more. However, when thecontent of S exceeds 0.05%, a low-melting point compound such as FeS maybe formed, which has a possibility of promoting high-temperature weldingcracks. Accordingly, the S content is not to be more than 0.05%.

[0065] The content of oxygen (O) is limited to 0.005% or less.

[0066] Where the content of O exceeds 0.005%, Ti forms Ti oxides inmolten steels, so that it cannot form TiN precipitates. Accordingly, itis undesirable for the O content to be more than 0.005%. Furthermore,inclusions such as coarse Fe oxides and Al oxides may be formed whichundesirably affect the toughness of the matrix.

[0067] In accordance with the present invention, the ratio of Ti/N islimited to a range of 1.2 to 2.5.

[0068] When the ratio of Ti/N is limited to a desired range as definedabove, there are two advantages as follows.

[0069] First, it is possible to increase the density of TiN precipitateswhile uniformly dispersing those TiN precipitates. That is, when thenitrogen content is increased under the condition in which the Ticontent is constant, all dissolved Ti atoms are easily coupled withnitrogen atoms in a continuous casing process (in the case of ahigh-nitrogen slab) or in a cooling process following a nitrogen zingtreatment (in the case of a low-nitrogen slab), so that fine TiNprecipitates are formed while being dispersed at an increased density.

[0070] Second, the solubility product of TiN representing thehigh-temperature stability of TiN precipitates is reduced, therebypreventing a re-dissolution of Ti. That is, Ti predominantly exhibits aproperty of coupling with N under a high-nitrogen environment, over adissolution property. Accordingly, TiN precipitates are stable at a hightemperature.

[0071] Therefore, the ratio of Ti/N is controlled to be 1.2 to 2.5 inaccordance with the present invention. When the Ti/N ratio is less than1.2, the amount of nitrogen dissolved in the matrix is increased,thereby degrading the toughness of the heat affected zone. On the otherhand, when the Ti/N ratio is more than 2.5, coarse TiN grains areformed. In this case, it is difficult to obtain a uniform dispersion ofTiN. Furthermore, the surplus Ti remaining without being precipitated inthe form of TiN is present in a dissolved state, so that it mayadversely affect the toughness of the heat affected zone.

[0072] The ratio of N/B is limited to a range of 10 to 40.

[0073] When the ratio of N/B is less than 10, BN serving to promote atransformation into polygonal ferrites at the grain boundaries of prioraustenite is precipitated in an insufficient amount in the coolingprocess following the welding process. On the other hand, when the N/Bratio exceeds 40, the effect of BN is saturated. In this case, anincrease in the amount of dissolved nitrogen occurs, thereby degradingthe toughness of the heat affected zone.

[0074] The ratio of Al/N is limited to a range of 2.5 to 7.

[0075] Where the ratio of Al/N is less than 2.5, AlN precipitates forcausing a transformation into acicular ferrites are dispersed at aninsufficient density. Furthermore, an increase in the amount ofdissolved nitrogen in the heat affected zone occurs, thereby possiblycausing formation of welding cracks. On the other hand, where the Al/Nratio exceeds 7, the effects obtained by controlling the Al/N ratio aresaturated.

[0076] The ratio of (Ti+2Al+4B)/N is limited to a range of 6.5 to 14.

[0077] Where the ratio of (Ti+2Al+4B)/N is less than 6.5, the grain sizeand density of TiN, AlN, BN, and VN precipitates are insufficient, sothat it is impossible to achieve suppression of the growth of prioraustenite grains in the heat affected zone, formation of fine polygonalferrite at grain boundaries, control of the amount of dissolvednitrogen, formation of acicular ferrite and polygonal ferrite withingrains, and control of structure fractions. On the other hand, when theratio of (Ti+2Al+4B)/N exceeds 14, the effects obtained by controllingthe ratio of (Ti+2Al+4B)/N are saturated. Where V is added, it ispreferable for the ratio of (Ti+2Al+4B+V)/N to range from 7 to 17.

[0078] The ratio of Mn/S is limited to a range of 220 to 400.

[0079] In accordance with the present invention, precipitates of MnS areformed at the boundaries between TiN precipitates and matrix.Accordingly, when these precipitates are heated to a high temperature,they are preferentially dissolved again in the matrix, therebyincreasing the re-dissolution temperature, as compared to TiNprecipitates dispersed alone, or delaying the time required forre-dissolution.

[0080] The ratio of Mn/S should be 220 or more in order to obtain anappropriate amount of complex precipitates of TiN and MnS for desiredcontrol of the growth of austenite grains in the heat affected zone.However, when the ratio of Mn/S exceeds 400, MnS precipitatessurrounding TiN precipitates are coarsened, so that the effects obtainedby controlling the ratio of Mn/S are saturated. Furthermore, an increasein the hardenability of the heat affected zone may occur, therebycausing a degradation in toughness while promoting formation ofhigh-temperature cracks in the welding metal.

[0081] In accordance with the present invention, V may also beselectively added to the above defined steel composition.

[0082] V is an element which is coupled with N to form VN, therebypromoting formation of ferrite in the heat affected zone. VN isprecipitated alone, or precipitated in TiN precipitates, so that itpromotes a ferrite transformation. Also, V is coupled with C, therebyforming a carbide, that is, VC. This VC serves to suppress growth offerrite grains after the ferrite transformation.

[0083] Thus, V further improves the toughness of the matrix and thetoughness of the heat affected zone. In accordance with the presentinvention, the content of V is preferably limited to a range of 0.01 to0.2%. Where the content of V is less than 0.01%, the amount ofprecipitated VN is insufficient to obtain an effect of promoting theferrite transformation in the heat affected zone. On the other hand,where the content of V exceeds 0.2%, both the toughness of the matrixand the toughness of the heat affected zone are degraded. In this case,an increase in welding hardenability occurs. For this reason, there is apossibility of formation of undesirable low-temperature welding cracks.

[0084] Where V is added, the ratio of V/N is preferably controlled to be0.3 to 9.

[0085] When the ratio of V/N is less than 0.3, it may be difficult tosecure an appropriate density and grain size of VN precipitatesdispersed at boundaries of complex precipitates of TiN and MnS for animprovement in the toughness of the heat affected zone. On the otherhand, when the ratio of V/N exceeds 9, the VN precipitates dispersed atthe boundaries of complex precipitates of TiN and MnS may be coarsened,thereby reducing the density of those VN precipitates. As a result, thefraction of ferrite effectively serving to improve the toughness of theheat affected zone may be reduced.

[0086] In order to further improve mechanical properties, the steelshaving the above defined composition may be added with one or moreelement selected from the group consisting of Ni, Cu, Nb, Mo, and Cr inaccordance with the present invention.

[0087] The content of Ni is preferably limited to a range of 0.1 to3.0%.

[0088] Ni is an element which is effective to improve the strength andtoughness of the matrix in accordance with a solid-solutionstrengthening. In order to obtain such an effect, the Ni content ispreferably 0.1% or more. However, when the Ni content exceeds 3.0%, anincrease in hardenability occurs, thereby degrading the toughness of theheat affected zone. Furthermore, there is a possibility of formation ofhigh-temperature cracks in both the heat affected zone and the matrix.

[0089] The content of copper (Cu) is limited to a range of 0.1 to 1.5%.

[0090] Cu is an element which is dissolved in the matrix, therebysolid-solution strengthening the matrix. That is, Cu is effective tosecure desired strength and toughness for the matrix. In order to obtainsuch an effect, Cu should be added in a content of 0.1% or more.However, when the Cu content exceeds 1.5%, the hardenability of the heataffected zone is increased, thereby causing a degradation in toughness.Furthermore, formation of high-temperature cracks at the heat affectedzone and welding metal is promoted. In particular, Cu is precipitated inthe form of CuS around Ti-based oxides, along with S, therebyinfluencing the formation of ferrites having an acicular or polygonalstructure effective to achieve an improvement in the toughness of theheat affected zone. Accordingly, it is preferred for the Cu content tobe 0.1 to 1.5%.

[0091] In the case where Cu and Ni are added together in presentinvention, it is preferable to limit the sum of the addition contents ofthem to a range of 3.5% or less. When the contents exceeds 3.5%, thetoughness of the heat affected zone and weldability could be aggravated.

[0092] The content of Nb is preferably limited to a range of 0.01 to0.10%.

[0093] Nb is an element which is effective to secure a desired strengthof the matrix. For such an effect, Nb is added in an amount of 0.01% ormore. However, when the content of Nb exceeds 0.1%, coarse NbC may beprecipitated alone, adversely affecting the toughness of the matrix.

[0094] The content of chromium (Cr) is preferably limited to a range of0.05 to 1.0%.

[0095] Cr serves to increase hardenability while improving strength. Ata Cr content of less than 0.05%, it is impossible to obtain desiredstrength. On the other hand, when the Cr content exceeds 1.0%, adegradation in toughness in both the matrix and the heat affected zoneoccurs.

[0096] The content of molybdenum (Mo) is preferably limited to a rangeof 0.05 to 1.0%.

[0097] Mo is an element which increases hardenability while improvingstrength. In order to secure desired strength, it is necessary to add Moin an amount of 0.05% or more. However, the upper limit of the Mocontent is determined to be 0.1%, similarly to Cr, in order to suppresshardening of the heat affected zone and formation of low-temperaturewelding cracks.

[0098] In accordance with the present invention, one or both of Ca andREM may also be added in order to suppress the growth of prior austenitegrains in a heating process.

[0099] Ca and REM serve to form an oxide exhibiting a superiorhigh-temperature stability, thereby suppressing the growth of prioraustenite grains in the matrix during a heating process while improvingthe toughness of the heat affected zone. Also, Ca has an effect ofcontrolling the shape of coarse MnS in a steel manufacturing process.For such effects, Ca is preferably added in an amount of 0.0005% ormore, whereas REM is preferably added in an amount of 0.005% or more.However, when the Ca content exceeds 0.005%, or the REM content exceeds0.05%, large-size inclusions and clusters are formed, thereby degradingthe cleanness of steels. For REM, one or more of Ce, La, Y, and Hf maybe used.

[0100] Now, the microstructure of the welding structural steel productaccording to the present invention will be described.

[0101] Preferably, the microstructure of the steel product according tothe present invention obtained after being subjected to a hot rollingprocess is a complex structure of ferrite and pearlite. Also, theferrite should have a grain size of 20 μm or less. Where ferrite grainshave a grain size of more than 20 μm, the prior austenite grains in theheat affected zone is rendered to have a grain size of 80 μm or morewhen a high heat input welding process is applied, thereby degrading thetoughness of the heat affected zone.

[0102] Where the fraction of ferrite in the complex structure of ferriteand pearlite is increased, the toughness and elongation of the matrixare correspondingly increased. Accordingly, the fraction of ferrite isdetermined to be 20% or more, and preferably 70% or more.

[0103] It is desirable that complex precipitates of TiN and MnS having agrain size of 0.01 to 0.1 μm are dispersed in the welding structuralsteel product (matrix) of the present invention at a density of1.0×10⁷/mm².

[0104] Where the precipitates have a grain size of less than 0.01 μm,they may be easily dissolved again in the matrix in a welding process,so that they cannot effectively suppress the growth of austenite grains.On the other hand, where the precipitates have a grain size of more than0.1 μm, they exhibit an insufficient pinning effect (suppression ofgrowth of grains) on austenite grains, and behave like as coarsenon-metallic inclusions, thereby adversely affecting mechanicalproperties. Where the density of the fine precipitates is less than1.0×10⁷/mm², it is difficult to control the critical austenite grainsize of the heat affected zone to be 80 μm or less where a weldingprocess using high input heat is applied.

[0105] Where the precipitates are uniformly dispersed, it is possible tomore effectively suppress the Ostwald ripening phenomenon causingcoarsening of precipitates. Accordingly, it is desirable to control TiNprecipitates to have a space of 0.5 μm.

[0106] [Method for Manufacturing Welding Structural Steel Products]

[0107] In accordance with the present invention, a steel slab having theabove defined composition is first prepared.

[0108] The steel slab of the present invention may be manufactured byconventionally processing, through a casting process, molten steeltreated by conventional refining and deoxidizing processes. However, thepresent invention is not limited to such processes.

[0109] In accordance with the present invention, molten steel isprimarily refined in a converter, and tapped into a ladle so that it maybe subjected to a “refining outside furnace” process as a secondaryrefining process. In the case of thick products such as weldingstructural steel products, it is desirable to perform a degassingtreatment (Ruhrstahi Hereaus (RH) process) after the “refining outsidefurnace” process. Typically, deoxidization is carried out between theprimary and secondary refining processes.

[0110] In the deoxidizing process, it is most desirable to add Ti underthe condition in which the amount of dissolved oxygen has beencontrolled not to be more than an appropriate level in accordance withthe present invention. This is because most of Ti is dissolved in themolten steel without forming any oxide. In this case, an element havinga deoxidizing effect higher than that of Ti is preferably added prior tothe addition of Ti.

[0111] This will be described in more detail. The amount of dissolvedoxygen greatly depends on an oxide production behavior. In the case ofdeoxidizing agents having a higher oxygen affinity, their rate ofcoupling with oxygen in molten steel is higher. Accordingly, where adeoxidation is carried out using an element having a deoxidizing effecthigher than that of Ti, prior to the addition of Ti, it is possible toprevent Ti from forming an oxide, as much as possible. Of course, adeoxidation may be carried out under the condition that Mn, Si, etc.belonging to the 5 elements of steel are added prior to the addition ofthe element having a deoxidizing effect higher than that of Ti, forexample, Al. After the deoxidation, a secondary deoxidation is carriedout using Al. In this case, there is an advantage in that it is possibleto reduce the amount of added deoxidizing agents. Respective deoxidizingeffects of deoxidizing agents are as follows:

Cr<Mn<Si<Ti<Al<REM<Zr<Ca≈Mg

[0112] As apparent from the above description, it is possible to controlthe amount of dissolved oxygen to be as low as possible by adding anelement having a deoxidizing effect higher than that of Ti, prior to theaddition of Ti, in accordance with the present invention. Preferably,the amount of dissolved oxygen is controlled to be 30 ppm or less. Whenthe amount of dissolved oxygen exceeds 30 ppm, Ti may be coupled withoxygen existing in the molten steel, thereby forming a Ti oxide. As aresult, the amount of dissolved Ti is reduced.

[0113] It is preferred that after the control of the dissolved oxygenamount, the addition of Ti be completed within 10 minutes under thecondition that the content of Ti ranges from 0.005% to 0.2%. This isbecause the amount of dissolved Ti may be reduced with the lapse of timedue to production of a Ti oxide after the addition of Ti.

[0114] In accordance with the present invention, the addition of Ti maybe carried out at any time before or after a vacuum degassing treatment.

[0115] In accordance with the present invention, a steel slab ismanufactured using the molten steel prepared as described above. Wherethe prepared molten steel is low-nitrogen steel (requiring anitrogenizing treatment), it is possible to carry out a continuouscasting process irrespective of its casting speed, that is, a lowcasting speed or a high casting speed. However, where the molten steelis high-nitrogen steel, it is desirable, in terms of an improvement inproductivity, to cast the molten steel at a low casting speed whilemaintaining a weak cooling condition in the secondary cooling zone,taking into consideration the fact that high-nitrogen steel has a highpossibility of formation of slab surface cracks.

[0116] Preferably, the casting speed of the continuous casting processis 1.1 m/min lower than a typical casting speed, that is, about 1.2m/min. More preferably, the casting speed is controlled to be about 0.9to 1.1 m/min. At a casting speed of less than 0.9 m/min, a degradationin productivity occurs even though there is an advantage in terms ofreduction of slab surface cracks. On the other hand, where the castingspeed is higher than 1.1 m/min, the possibility of formation of slabsurface cracks is increased. Even in the case of low-nitrogen steel, itis possible to obtain a better internal quality when the steel is castat a low speed of 0.9 to 1.2 m/min.

[0117] Meanwhile, it is desirable to control the cooling condition atthe secondary cooling zone because the cooling condition influences thefineness and uniform dispersion of TiN precipitates.

[0118] For high-nitrogen molten steel, the water spray amount in thesecondary cooling zone is determined to be 0.3 to 0.35 l/kg for weakcooling. When the water spray amount is less than 0.3 l/kg, coarseningof TiN precipitates occurs. As a result, it may be difficult to controlthe grain size and density of TiN precipitates in order to obtaindesired effects according to the present invention. On the other hand,when the water spray amount is more than 0.35 l/kg, the frequency offormation of TiN precipitates is too low so that it is difficult tocontrol the grain size and density of TiN precipitates in order toobtain desired effects according to the present invention.

[0119] Thereafter, the steel slab prepared as described above is heatedin accordance with the present invention.

[0120] In the case of a high-nitrogen steel slab having a nitrogencontent of 0.008 to 0.030%, it is heated at a temperature of 1,100 to1,250° C. for 60 to 180 minutes. When the slab heating temperature isless than 1,100° C., it is difficult to secure the grain sizes anddensities of precipitates of MnS and complex precipitates of TiN and MnSappropriate to obtain desired effects according to the presentinvention. On the other hand, when the slab heating temperature is morethan 1,250° C., the grain size and density of complex precipitates ofTiN and MnS are saturated. Also, austenite grains are grown during theheating process. As a result, the austenite grains, which influencerecrystallization to be performed in a subsequent rolling process, areexcessively coarsened, so that they exhibit a reduced effect of finingferrite, thereby degrading the mechanical properties of the final steelproduct.

[0121] Meanwhile, where the slab heating time is less than 60 minutes,solidification segregation is reduced. Also, the given time isinsufficient to allow complex precipitates of TiN and MnS to bedispersed. When the heating time exceeds 180 minutes, the effectsobtained by the heating process are saturated. In this case, there is anincrease in the manufacturing costs. Furthermore, growth of austenitegrains occurs in the slab, adversely affecting the subsequent rollingprocess.

[0122] For a low-nitrogen steel slab containing nitrogen in an amount of0.005%, a nitrogenizing treatment is carried out in a slab heatingfurnace in accordance with the present invention so as to obtain ahigh-nitrogen steel slab while adjusting the ratio between Ti and N.

[0123] In accordance with the present invention, the high-nitrogen slabis heated at a temperature of 1,000 to 1,250° C. for 60 to 180 minutesfor a nitrogenizing treatment thereof, in order to control the nitrogenconcentration of the slab to be preferably 0.008 to 0.03%. In order tosecure an appropriate amount of TiN precipitates in the slab, thenitrogen content should be 0.008% or more. However, when the nitrogencontent exceeds 0.03%, nitrogen may be diffused in the slab, therebycausing the amount of nitrogen at the surface of the slab to be morethan the amount of nitrogen precipitated in the form of fine TiNprecipitates. As a result, the slab is hardened at its surface, therebyadversely affecting the subsequent rolling process.

[0124] When the heating temperature of the slab is less than 1,000° C.,nitrogen cannot be sufficiently diffused, thereby causing fine TiNprecipitates to have a low density. Although it is possible to increasethe density of TiN precipitates by increasing the heating time, thiswould increase the manufacturing costs. On the other hand, when theheating temperature is more than 1,250° C., growth of austenite grainsoccurs in the slab during the heating process, adversely affecting therecrystallization to be performed in the subsequent rolling process.Where the slab heating time is less than 60 minutes, it is impossible toobtain a desired nitrogenizing effect. On the other hand, where the slabheating time is more than 180 minutes, the manufacturing costsincreases. Furthermore, growth of austenite grains occurs in the slab,adversely affecting the subsequent rolling process.

[0125] Preferably, the nitrogenizing treatment is performed to control,in the slab, the ratio of Ti/N to be 1.2 to 2.5, the ratio of N/B to be10 to 40, the ratio of Al/N to be 2.5 to 7, the ratio of (Ti+2Al+4B)/Nto be 6.5 to 14, the ratio of V/N to be 0.3 to 9, and the ratio of(Ti+2Al+4B+V)/N to be 7 to 17.

[0126] Thereafter, the heated steel slab is hot-rolled in an austeniterecrystallization temperature range at a thickness reduction rate of 40%or more. The austenite recrystallization temperature range depends onthe composition of the steel, and a previous thickness reduction rate.In accordance with the present invention, the austeniterecrystallization temperature range is determined to be about 850 to1,050° C., taking into consideration a typical thickness reduction rate,along with the steel composition of the present invention.

[0127] Where the hot rolling temperature is less than 850° C., thestructure is changed into elongated austenite in the rolling processbecause the hot rolling temperature is within a non-crystallizationtemperature range. For this reason, it is difficult to secure fineferrite in a subsequent cooling process. On the other hand, where thehot rolling temperature is more than 1,050° C., grains of recrystallizedaustenite formed in accordance with recrystallization are grown, so thatthey are coarsened. As a result, it is difficult to secure fine ferritegrains in the cooling process. Also, when the accumulated or singlethickness reduction rate in the rolling process is less then 40%, thereare insufficient sites for formation of ferrite nuclei within austenitegrains. As a result, it is impossible to obtain an effect ofsufficiently fining ferrite grains in accordance with recrystallizationof austenite. Furthermore, there is an adverse affect on the behavior ofprecipitates advantageously influencing the toughness of the heataffected zone in a welding process.

[0128] The rolled steel slab is then cooled to a temperature ranging±10° C. from a ferrite transformation finish temperature at a rate of 1°C./min. Preferably, the rolled steel slab is cooled to the ferritetransformation finish temperature at a rate of 1° C./min, and thencooled in air.

[0129] Of course, there is no problem associated with fining of ferriteeven when the rolled steel slab is cooled to normal temperature at arate of 1° C./min. However, this is undesirable because it isuneconomical. Although the rolled steel slab is cooled to a temperatureranging ±10° C. from the ferrite transformation finish temperature at arate of 1° C./min, it is possible to prevent growth of ferrite grains.When the cooling rate is less than 1° C./min, growth of recrystallizedfine ferrite grains occurs. In this case, it is difficult to secure aferrite grain size of 20 μm or less.

[0130] It is possible to obtain a steel product having a complexstructure of ferrite and pearlite as its microstructure while exhibitinga superior heat affected zone toughness by controlling deoxidizing andcasting conditions while regulating content ratios of elements, inparticular, the ratio of Ti/N. Also, it is possible to effectivelymanufacture a steel product in which complex precipitates of TiN and MnShaving a grain size of 0.01 to 0.1 μm are precipitated at a density of1.0×10⁷/mm² or more while having a space of 0.5 μm or less.

[0131] Meanwhile, slabs can be manufactured using a continuous castingprocess or a mold casting process as a steel casting process. Where ahigh cooling rate is used, it is easy to finely disperse precipitates.Accordingly, it is desirable to use a continuous casting process. Forthe same reason, it is advantageous for the slab to have a smallthickness. As the hot rolling process for such a slab, a hot chargerolling process or a direct rolling process may be used. Also, varioustechniques such as known control rolling processes and controlledcooling processes may be employed. In order to improve the mechanicalproperties of hot-rolled plates manufactured in accordance with thepresent invention, a heat treatment may be applied. It should be notedthat although such known techniques are applied to the presentinvention, such an application is made within the scope of the presentinvention.

[0132] [Welded Structures]

[0133] The present invention also relates to a welded structuremanufactured using the above described welding structural steel product.Therefore, included in the present invention are welded structuresmanufactured using a welding structural steel product having the abovedefined composition according to the present invention, a microstructurecorresponding to a complex structure of ferrite and pearlite having agrain size of about 20 μm or less, or complex precipitates of TiN andMnS having a grain size of 0.01 to 0.1 μm while being dispersed at adensity of 1.0×10⁷/mm² or more and with a spacing of 0.5 μm or less.

[0134] Where a high heat input welding process is applied to the abovedescribed welding structural steel product, prior austenite having agrain size of 80 μm or less is formed. Where the grain size of the prioraustenite is more than 80 μm, an increase in hardenability occurs,thereby causing easy formation of a low-temperature structure(martensite or upper bainite). Furthermore, although ferrites havingdifferent nucleus forming sites are formed at grain boundaries ofaustenite, they are merged together when growth of grains occurs,thereby causing an adverse effect on toughness.

[0135] When the steel product is quenched in accordance with anapplication of a high heat input welding process thereto, themicrostructure of the heat affected zone includes ferrite having a grainsize of 20 μm or less at a volume fraction of 70% or more. Where thegrain size of the ferrite is more than 20 μm, the fraction of side plateor allotriomorphs ferrite adversely affecting the toughness of the heataffected zone increases. In order to achieve an improvement intoughness, it is desirable to control the volume fraction of ferrite tobe 70% or more. When the ferrite of the present invention hascharacteristics of polygonal ferrite or acicular ferrite, an improvementin toughness is expected. In accordance with the present invention, thiscan be achieved by forming BN and Fe-based carbide boride.

[0136] When a high heat input welding process is applied to the weldingstructural steel product (matrix), prior austenite having a grain sizeof 80 μm or less is formed at the heat affected zone. In accordance witha subsequent quenching process, the microstructure of the heat affectedzone includes ferrite having a grain size of 20 μm or less at a volumefraction of 70% or more.

[0137] Where a welding process using a heat input of 100 kJ/cm or lessis applied to the welding structural steel product of the presentinvention (in the case “Δt₈₀₀₋₅₀₀=60 seconds” in Table 5), the toughnessdifference between the matrix and the heat affected zone is within arange of ±30 J. In the case of a welding process using a high heat inputof 100 to 250 kJ/cm or more (“Δ₈₀₀₋₅₀₀=120 seconds” in Table 5), thetoughness difference between the matrix and the heat affected zone iswithin a range of ±40 J. Also, in the case of a welding process using ahigh heat input of 250 kJ/cm or more (“Δt₈₀₀₋₅₀₀=180 seconds” in Table5), the toughness difference between the matrix and the heat affectedzone is within a range of 0 to 100 J. Such results can be seen from thefollowing examples.

EXAMPLES

[0138] Hereinafter, the present invention will be described inconjunction with various examples. These examples are made only forillustrative purposes, and the present invention is not to be construedas being limited to those examples.

Example 1

[0139] Each of steel products having different steel compositions ofTable 1 was melted in a converter. The resultant molten steel wassubjected to a continuous casting process after being refined under thecondition of Table 2, thereby manufacturing a slab. The slab was thenhot rolled under the condition of Table 4, thereby manufacturing ahot-rolled plate. Table 3 describes content ratios of alloying elementsin each steel product. TABLE 1 Chemical Composition (wt %) B N C Si Mn PS Al Ti (ppm) (ppm) W Present 0.12 0.13 1.54 0.006 0.005 0.04 0.014 7120 0.005 Steel 1 Present 0.07 0.12 1.71 0.006 0.006 0.07 0.05 10 2800.002 Steel 2 Present 0.14 0.10 2.01 0.006 0.008 0.06 0.015 3 110 0.003Steel 3 Present 0.10 0.12 1.80 0.006 0.007 0.02 0.02 5 80 0.001 Steel 4Present 0.08 0.15 2.1 0.006 0.006 0.09 0.05 15 300 0.002 Steel 5 Present0.10 0.14 2.0 0.007 0.005 0.025 0.02 10 100 0.004 Steel 6 Present 0.130.14 1.6 0.007 0.007 0.04 0.015 8 115 0.15  Steel 7 Present 0.11 0.151.52 0.007 0.006 0.06 0.018 10 120 0.001 Steel 8 Present 0.13 0.21 1.420.007 0.005 0.025 0.02 4 90 0.002 Steel 9 Present 0.07 0.16 2.2 0.0080.010 0.045 0.025 6 100 0.05  Steel 10 Present 0.11 0.21 1.48 0.0070.006 0.047 0.019 11 130 0.01  Steel 11 Conventional 0.05 0.13 1.310.002 0.006 0.0014 0.009 1.6 22 — Steel 1 Conventional 0.05 0.11 1.340.002 0.003 0.0036 0.012 0.5 48 — Steel 2 Conventional 0.13 0.24 1.440.0012 0.003 0.0044 0.010 1.2 127 — Steel 3 Conventional 0.06 0.18 1.350.008 0.002 0.0027 0.013 8 32 — Steel 4 Conventional 0.06 0.18 0.880.006 0.002 0.0021 0.013 5 20 — Steel 5 Conventional 0.13 0.27 0.980.005 0.001 0.001 0.009 11 28 — Steel 6 Conventional 0.13 0.24 1.440.004 0.002 0.02 0.008 8 79 — Steel 7 Conventional 0.07 0.14 1.52 0.0040.002 0.002 0.007 4 57 — Steel 8 Conventional 0.06 0.25 1.31 0.008 0.0020.019 0.007 10 91 — Steel 9 Conventional 0.09 0.26 0.86 0.009 0.0030.046 0.008 15 142 — Steel 10 Conventional 0.14 0.44 1.35 0.012 0.0120.030 0.049 7 89 — Steel 11 Chemical Composition (wt %) O Cu Ni Cr Mo NbV Ca REM (ppm) Present 0.1 — — — — 0.01 — — 11 Steel 1 Present — 0.2 — —— 0.01 — — 12 Steel 2 Present — — — — — 0.02 — — 10 Steel 3 Present 0.1— — — — 0.5  — —  9 Steel 4 Present — — 0.1 — — 0.05 — — 12 Steel 5Present — — — 0.1 — 0.09 — —  9 Steel 6 Present 0.1 — — — — 0.02 — — 11Steel 7 Present — — — —  0.015 0.01 — — 10 Steel 8 Present — — 0.1 — —0.02 0.001 — 12 Steel 9 Present — 0.3 — — 0.01 0.02 — 0.01 11 Steel 10Present — 0.1 — — — — — — 15 Steel 11 Conventional — — — — — — — — 22Steel 1 Conventional — — — — — — — — 32 Steel 2 Conventional 0.3 — — —0.05 — — — 138  Steel 3 Conventional — —  0.14 0.15 —  0.028 — — 25Steel 4 Conventional  0.75  0.58  0.24 0.14  0.015  0.037 — — 27 Steel 5Conventional  0.35  1.15  0.53 0.49  0.001  0.045 — — 25 Steel 6Conventional 0.3 — — —  0.036 — — — Steel 7 Conventional  0.32  0.35 — — 0.013 — — — — Steel 8 Conventional — —  0.21  0.19  0.025  0.035 — — —Steel 9 Conventional —  1.09  0.51  0.36  0.021  0.021 — — — Steel 10Conventional — — — — —  0.069 — — — Steel 11

[0140] TABLE 2 Dissolved Oxygen Amount of Ti Water Primary Amount afterAdded after Casting Spray Steel Deoxidation Addition of DeoxidationSpeed Amount Products Sample Order Al (ppm) (%) (m/min) (l/kg) PS* 1 PS1 Mn → Si 19 0.014 1.1 0.32 PS 2 Mn → Si 18 0.014 1.1 0.32 PS 3 Mn → Si18 0.014 1.1 0.32 CS 1 Mn → Si 32 0.014 1.1 0.32 CS 2 Mn → Si 58 0.0141.1 0.32 PS* 2 PS 4 Mn → Si 16 0.05 1.0 0.35 PS* 3 PS 5 Mn → Si 15 0.0151.0 0.35 PS* 4 PS 6 Mn → Si 15 0.02 1.0 0.35 PS* 5 PS 7 Mn → Si 12 0.051.2 0.30 PS* 6 PS 8 Mn → Si 17 0.02 1.2 0.30 PS* 7 PS 9 Mn → Si 18 0.0151.1 0.32 PS* 8 PS 10 Mn → Si 14 0.018 1.1 0.32 PS* 9 PS 11 Mn → Si 190.02 1.1 0.32 PS* 10 PS 12 Mn → Si 22 0.025 1.0 0.35 PS* 11 PS 13 Mn →Si 20 0.019 1.0 0.35

[0141] TABLE 3 Content Ratios of Alloying Elements Steel Product Mn/STi/N N/B Al/N V/N (Ti + 2Al + 4B + V)/N Present 308 1.2 17.1 3.3 0.8 8.9Sample 1 Present 308 1.2 17.1 3.3 0.8 8.9 Sample 2 Present 308 1.2 17.13.3 0.8 8.9 Sample 3 Present 285 1.8 28.0 2.5 0.4 7.3 Sample 4 Present251 1.4 36.7 5.5 1.8 14.2 Sample 5 Present 257 2.5 16.0 2.5 6.3 14.0Sample 6 Present 350 1.7 20.0 3.0 1.7 9.5 Sample 7 Present 400 2.0 10.02.5 9.0 16.4 Sample 8 Present 229 1.3 14.4 3.5 1.7 10.3 Sample 9 Present253 1.5 12.0 5.0 0.8 12.7 Sample 10 Present 284 2.2 22.5 2.8 2.2 10.2Sample 11 Present 220 2.5 16.7 4.5 2.0 13.7 Sample 12 Present 247 1.511.8 3.6 — 9.0 Sample 13 Conventional 218 4.1 13.8 0.6 — 5.7 Steel 1Conventional 447 2.5 96.0 0.8 — 4.0 Steel 2 Conventional 480 0.8 105.80.4 — 1.5 Steel 3 Conventional 657 4.1 4.0 0.8 8.8 15.5 Steel 4Conventional 440 6.5 4.0 1.1 18.5  28.1 Steel 5 Conventional 980 3.2 2.60.4 16.1  21.6 Steel 6 Conventional 720 1.0 9.9 2.5 — 6.5 Steel 7Conventional 760 1.2 14.3 0.4 — 2.2 Steel 8 Conventional 655 0.8 9.1 2.13.9 9.2 Steel 9 Conventional 287 0.6 9.5 3.2 1.5 8.9 Steel 10Conventional 113 5.5 12.7 3.4 7.8 20.3 Steel 11

[0142] TABLE 4 Heating TRR(%)/ Cooling Heating Time Rolling StartRolling End ATRR Rate Steel Products Samples Temp. (° C.) (min) Temp. (°C.) Time (° C.) (%) (° C./min) Present Sample 2 Present 1150 170 1030850 65/80 5 Example 1 Present 1200 130 1040 850 65/80 5 Example 2Present 1240  90 1040 850 65/80 5 Example 3 Comparative 1050  60 1040850 65/80 5 Example 1 Comparative 1300 250 1035 850 65/80 5 Example 2Present Sample 1 Present 1200 130 1020 840 65/80 6 Example 4 PresentSample 3 Present 1200 130 1040 850 65/80 6 Example 5 ComparativeComparative 1210 120 1030 860 65/80 0.01 Sample 1 Example 3 ComparativeComparative 1210 120 1030 860 65/80 35 Sample 2 Example 4 Present Sample4 Present 1180 150 1020 860 60/80 5 Example 6 Present Sample 5 Present1190 140 1010 850 60/80 5 Example 7 Present Sample 6 Present 1220 1101010 840 60/75 6 Example 8 Present Sample 7 Present 1220 110 1020 84060/75 7 Example 9 Present Sample 8 Present 1210 120 1010 850 60/75 7Example 10 Present Sample 9 Present 1220 110 1000 840 55/70 5 Example 11Present Sample 10 Present 1210 120 1010 830 55/70 6 Example 12 PresentSample 11 Present 1230 100 1000 850 55/70 5 Example 13 Present Sample 12Present 1220 110 1020 840 55/70 5 Example 14 Present Sample 13 Present1210 130 1020 840 65/75 5 Example 15 Conventional Steel 11 1200 — Ar₃960 80 Naturally or more Cooled

[0143] Test pieces were sampled from the hot-rolled products. Thesampling was performed at the central portion of each hot-rolled productin a thickness direction. In particular, test pieces for a tensile testwere sampled in a rolling direction, whereas test pieces for a Charpyimpact test were sampled in a direction perpendicular to the rollingdirection. Using steel test pieces sampled as described above,characteristics of precipitates in each steel product (matrix), andmechanical properties of the steel product were measured. The measuredresults are described in Table 5. Also, the microstructure and impacttoughness of the heat affected zone were measured. The measured resultsare described in Table 6.

[0144] These measurements were carried out as follows.

[0145] For tensile test pieces, test pieces of KS Standard No. 4 (KS B0801) were used. The tensile test was carried out at a cross heat speedof 5 mm/min. On the other hand, impact test pieces were prepared, basedon the test piece of KS Standard No. 3 (KS B 0809). For the impact testpieces, notches were machined at a side surface (L-T) in a rollingdirection in the case of the matrix while being machined in a weldingline direction in the case of the welding material. In order to inspectthe size of austenite grains at a maximum heating temperature of theheat affected zone, each test piece was heated to a maximum heatingtemperature of 1,200 to 1,400° C. at a heating rate of 140° C./sec usinga reproducible welding simulator, and then quenched using He gas afterbeing maintained for one second. After the quenched test piece waspolished and eroded, the grain size of austenite in the resultant testpiece at a maximum heating temperature condition was measured inaccordance with a KS Standard (KS D 0205).

[0146] The microstructure obtained after the cooling process, and thegrain sizes, densities, and spacing of precipitates and oxides seriouslyinfluencing the toughness of the heat affected zone were measured inaccordance with a point counting scheme using an image analyzer and anelectronic microscope. The measurement was carried out for a test areaof 100 mm². The impact toughness of the heat affected zone in each testpiece was evaluated by subjecting the test piece to welding conditionscorresponding to welding heat inputs of about 80 kJ/cm, 150 kJ/cm, and250 kJ/cm, that is, welding cycles involving heating at a maximumheating temperature of 1,400° C., and cooling for 60 seconds, 120seconds, and 180 seconds, respectively, polishing the surface of thetest piece, machining the test piece for an impact test, and thenconducting a Charpy impact test for the test piece at a temperature of−40° C. TABLE 5 Characteristics of Characteristics of Matrix Structureand Precipitates of Mechanical Properties of Matrix TiN + MnS Volume−40° C. Density Mean Thick- Yield Tensile Elonga- Fraction Impact(number/ Size Spacing ness Strength Strength tion FGS of FerriteToughness Sample mm²) (μm) (μm) (mm) (MPa) (MPa) (%) (μm) (%) (J) PE 12.4 × 10⁸ 0.016 0.25 25 394 553 38 11 82 358 PE 2 3.2 × 10⁸ 0.017 0.2425 395 551 39 9 83 362 PE 3 2.5 × 10⁸ 0.012 0.26 25 396 550 39 10 83 357CE 1 2.3 × 10⁶ 0.174 1.6 25 393 554 26 16 70 206 CE 2 3.4 × 10⁶ 0.1651.8 25 792 860 17 17 21 45 PE 4 3.2 × 10⁸ 0.025 0.32 30 396 558 38 11 83349 PE 5 2.6 × 10⁸ 0.013 0.34 30 396 562 38 10 83 354 CE 3 1.3 × 10⁶0.182 1.2 30 384 564 30 18 73 220 CE 4 4.3 × 10⁶ 0.177 1.4 30 392 582 2917 74 208 PE 6 3.3 × 10⁸ 0.026 0.35 30 390 563 38 10 82 364 PB 7 4.6 ×10⁸ 0.024 0.32 35 390 564 39 10 85 360 PE 8 4.3 × 10⁸ 0.014 0.40 35 392542 36 11 82 365 PE 9 5.6 × 10⁸ 0.028 0.29 35 391 536 37 10 84 359 PE 105.2 × 10⁸ 0.021 0.28 35 394 566 36 10 83 375 PE 11 3.7 × 10⁸ 0.029 0.2540 390 566 37 12 83 364 PE 12 3.2 × 10⁸ 0.025 0.31 40 396 542 38 11 85356 PE 13 3.3 × 10⁸ 0.042 0.34 40 406 564 38 12 82 348 PE 14 3.6 × 10⁸0032 0.28 40 387 550 37 10 83 349 PE 15 4.2 × 10⁸ 0.018 0.26 30 389 54939 9 86 368 CS 1 35 406 436 CS 2 35 405 441 CS 3 25 629 681 CS 4Precipitates of MgO—TiN 40 472 609 3.03 × 10⁶/mm² CS 5 Precipitates ofMgO—TiN 40 494 622 4.07 × 10⁶/mm² CS 6 Precipitates of MgO—TiN 50 812912 2.80 × 10⁶/mm² CS 7 25 629 681 CS 8 50 504 601 CS 9 60 526 648 CS 1060 760 829 CS 11 50 401 514

[0147] Referring to Table 5, it can be seen that the density ofprecipitates (complex precipitates of TiN and MnS) in each hot-rolledproduct manufactured in accordance with the present invention is1.0×10⁸/mm² or more, whereas the density of precipitates in eachconventional product is 4.07×10⁵/mm² or less. That is, the product ofthe present invention is formed with precipitates having a very smallgrain size while being dispersed at a considerably increased density.

[0148] The products of the present invention have a matrix structurehaving fine ferrite with a grain size of about 8 μm or less at a highfraction of 87% or more. TABLE 6 Microstructure Mechanical ReproducibleHeat Affected Zone of Heat Affected Properties of Impact Toughness (J)at −40° C. Grain Size of Zone with Heat Welded Zone (Maximum HeatingTemp. 1,400° C.) Austenite in Input of 100 kJ/cm Δ t₈₀₀₋₅₀₀ = Δ t₈₀₀₋₅₀₀= Δ t₈₀₀₋₅₀₀ = Heat Affected Volume Mean 180 sec 120 sec 180 sec Zone(μm) Fraction Grain Yield Tensile Impact Transition Impact Transition1,200 1,300 1400 of Ferrite Size of Strength Strength Toughness Temp.Toughness Temp. Sample (° C.) (° C.) (° C.) (%) Ferrite (μm) (kg/mm²)(kg/mm²) (J) (° C.) (J) (° C.) PE 1 23 33 56 73 16 370 −74 330 −67 294−62 PE 2 22 34 55 76 15 383 −76 353 −69 301 −63 PE 3 23 32 56 74 17 365−72 331 −67 298 −63 CE 1 54 84 182 36 32 126 −43 47 −34 26 −27 CE 2 6591 198 37 35 104 −40 35 −32 18 −26 PE 4 25 37 65 75 18 353 −71 325 −68287 −64 PE 5 26 40 57 74 16 362 −71 333 −67 296 −61 CE 3 48 78 220 58 22182 −44 87 −36 36 −28 CE 4 56 82 254 52 26 176 −44 79 −35 32 −29 PE 6 2531 53 76 17 386 −73 353 −69 305 −62 PE 7 24 34 55 74 18 367 −71 338 −67293 −63 PE 8 27 36 53 73 14 364 −71 334 −67 294 −61 PE 9 24 36 52 74 17367 −72 335 −67 285 −62 PE 10 22 35 53 73 18 385 −72 345 −66 294 −61 PE11 26 34 64 74 16 358 −71 324 −68 285 −63 PE 12 27 38 64 74 18 355 −71324 −67 284 −62 PE 13 24 32 54 75 16 367 −72 336 −68 285 −63 PE 14 25 3158 72 17 365 −72 330 −68 280 −63 PE 15 24 32 54 76 14 368 −72 345 −68286 −63 CS 1 187 −51 CS 2 156 −48 CS 3 148 −50 CS 4 230 93 143 −48132(0° C.) CS 5 180 87 132 −45 129(0° C.) CS 6 250 47 153 −43  60(0° C.)CS 7 141 −54 −61 CS 8 156 −59 −48 CS 9 145 −54 −42 CS 10 138 −57 −45 CS11 141 −43 219(0° C.)

[0149] Referring to Table 6, it can be seen that the size of austenitegrains under a maximum heating temperature condition of 1,400° C., as inthe heat affected zone, is within a range of 52 to 65 μm in the case ofthe present invention, whereas the austenite grains in the conventionalproducts are very coarse to have a grain size of about 180 μm. Thus, thesteel products of the present invention have a superior effect ofsuppressing the growth of austenite grains at the heat affected zone ina welding process. Where a welding process using a heat input of 100kJ/cm is applied, the steel products of the present invention have aferrite fraction of about 70% or more.

[0150] Under a high heat input welding condition in which a welding heatinput is 250 kJ/cm (the time taken for cooling from 800° C. to 500° C.is 180 seconds), the products of the present invention exhibit asuperior toughness value of about 280 J or more as a heat affected zoneimpact toughness at −40° C. while exhibiting about −60° C. as atransition temperature. That is, the products of the present inventionexhibit a superior heat affected zone impact toughness under a high heatinput welding condition.

[0151] Under the same high heat input welding condition, theconventional steel products exhibit a toughness value of about 200 J asa heat affected zone impact toughness at 0° C. while exhibiting about−60° C. as a transition temperature.

Example 2 Nitrogenizing Treatment

[0152] Each of steel products having different steel compositions ofTable 7 was melted in a converter. The resultant molten steel wassubjected to a continuous casting process after being deoxidized whilebeing subsequently added with Ti, thereby manufacturing a slab.

[0153] The slab was then hot rolled under the condition of Table 9,thereby manufacturing a hot-rolled plate. Table 10 describes contentratios of alloying elements in each steel product. TABLE 7 ChemicalComposition (wt %) B N C Si Mn P S Al Ti (ppm) (ppm) W Present 0.12 0.131.54 0.006 0.005 0.04 0.014 7 40 0.005 Steel 1 Present 0.07 0.12 1.500.006 0.005 0.07 0.05 10 43 0.002 Steel 2 Present 0.14 0.10 1.55 0.0060.007 0.06 0.015 3 41 0.003 Steel 3 Present 0.10 0.12 1.48 0.006 0.0050.02 0.02 5 40 0.001 Steel 4 Present 0.08 0.15 1.52 0.006 0.004 0.090.05 15 43 0.002 Steel 5 Present 0.10 0.14 1.50 0.007 0.005 0.025 0.0210 40 0.004 Steel 6 Present 0.13 0.14 1.59 0.007 0.007 0.04 0.015 8 450.15  Steel 7 Present 0.11 0.15 1.54 0.007 0.007 0.06 0.018 10 42 0.001Steel 8 Present 0.13 0.21 1.50 0.007 0.005 0.025 0.02 4 40 0.002 Steel 9Present 0.07 0.16 1.45 0.008 0.006 0.045 0.025 6 41 0.05  Steel 10Present 0.09 0.21 1.47 0.006 0.004 0.047 0.019 11 42 0.01  Steel 11Conventional 0.05 0.13 1.31 0.002 0.006 0.0014 0.009 1.6 22 — Steel 1Conventional 0.05 0.11 1.34 0.002 0.003 0.0036 0.012 0.5 48 — Steel 2Conventional 0.13 0.24 1.44 0.0012 0.003 0.0044 0.010 1.2 127 — Steel 3Conventional 0.06 0.18 1.35 0.008 0.002 0.0027 0.013 8 32 — Steel 4Conventional 0.06 0.18 0.88 0.006 0.002 0.0021 0.013 5 20 — Steel 5Conventional 0.13 0.27 0.98 0.005 0.001 0.001 0.009 11 28 — Steel 6Conventional 0.13 0.24 1.44 0.004 0.002 0.02 0.008 8 79 — Steel 7Conventional 0.07 0.14 1.52 0.004 0.002 0.002 0.007 4 57 — Steel 8Conventional 0.06 0.25 1.31 0.008 0.002 0.019 0.007 10 91 — Steel 9Conventional 0.09 0.26 0.86 0.009 0.003 0.046 0.008 15 142 — Steel 10Conventional 0.14 0.44 1.35 0.012 0.012 0.030 0.049 7 89 — Steel 11Chemical Composition (wt %) O Cu Ni Cr Mo Nb V Ca REM (ppm) Present 0.2— — — — 0.01 — — 11 Steel 1 Present 0.1 0.2 — — — 0.01 — — 12 Steel 2Present 0.1 — — — — 0.02 — — 10 Steel 3 Present 0.3 — — — — 0.5  — — 9Steel 4 Present 0.1 — 0.1 — — 0.05 — — 12 Steel 5 Present  0.45 — — 0.1— 0.09 — — 9 Steel 6 Present 0.1 — — — — 0.02 — — 11 Steel 7 Present 0.3— — — 0.015 0.01 — — 10 Steel 8 Present  0.21 — 0.1 — — 0.02 0.001 — 12Steel 9 Present 0.1 0.3 — — 0.01  0.02 — 0.01 8 Steel 10 Present 0.2 0.1— — — — — — 14 Steel 11 Conventional — — — — — — — — 22 Steel 1Conventional — — — — — — — — 32 Steel 2 Conventional 0.3 — — — 0.05  — —— 138 Steel 3 Conventional — —  0.14 0.15 —  0.028 — — 27 Steel 4Conventional  0.75  0.68  0.24 0.14 0.015  0.037 — — 25 Steel 5Conventional  0.35  1.15  0.53 0.49 0.001  0.045 — — — Steel 6Conventional 0.3 — — — 0.036 — — — Steel 7 Conventional  0.32  0.35 — —0.013 — — — — Steel 8 Conventional — —  0.21 0.19 0.025  0.035 — — —Steel 9 Conventional —  1.09  0.51 0.36 0.021  0.021 — — — Steel 10Conventional — — — — —  0.069 — — — Steel 11

[0154] TABLE 8 Dissolved Oxygen Amount of Ti Water Primary Amount afterAdded after Casting Spray Steel Deoxidation Addition of DeoxidationSpeed Amount Products Sample Order Al (ppm) (%) (m/min) (l/kg) PS* 1 PS1 Mn → Si 19 0.014 1.1 0.32 PS 2 Mn → Si 18 0.014 1.1 0.32 PS 3 Mn → Si18 0.014 1.1 0.32 CS 1 Mn → Si 32 0.014 1.1 0.32 CS 2 Mn → Si 58 0.0141.1 0.32 PS* 2 PS 4 Mn → Si 16 0.05 1.0 0.35 PS* 3 PS 5 Mn → Si 15 0.0151.0 0.35 PS* 4 PS 6 Mn → Si 15 0.02 1.0 0.35 PS* 5 PS 7 Mn → Si 12 0.051.2 0.30 PS* 6 PS 8 Mn → Si 17 0.02 1.2 0.30 PS* 7 PS 9 Mn → Si 18 0.0151.1 0.32 PS* 8 PS 10 Mn → Si 14 0.018 1.1 0.32 PS* 9 PS 11 Mn → Si 190.02 1.1 0.32 PS* 10 PS 12 Mn → Si 22 0.025 1.0 0.35 PS* 11 PS 13 Mn →Si 20 0.019 1.0 0.35

[0155] TABLE 9 Rolling Rolling TRR(%)/ATRR Nitrogen HeatingNitrogenizing Heating Start End (%) in Cooling Content of Steel Temp.Atmosphere Time Temp. Temp. Recrystallization Rate Matrix ProductsSample (° C.) (l/min) (min) (° C.) (° C.) Range (° C./min) (ppm) PS* 1PS 1 1220 350 160 1030 830 55/75 5 105 PS 2 1190 610 120 1020 830 55/755 115 PS 3 1150 780 100 1020 830 55/75 5 120 CS 1 1050 220  60 1020 84055/75 5 72 CS 2 1300 950 180 1020 840 55/75 5 316 PS* 2 PS 4 1180 780110 1010 830 55/75 6 275 PS* 3 PS 5 1200 600 100 1040 850 55/75 7 112PS* 4 PS 6 1170 620 130 1030 840 55/75 7 80 PS* 5 PS 7 1190 780 100 1020830 55/75 6 300 PS* 6 PS 8 1200 620 110 1030 830 55/75 6 100 PS* 7 PS 91150 750 160 1040 830 60/70 6 115 PS* 8 PS 10 1180 630 110 1040 85060/70 5 120 PS* 9 PS 11 1200 520 100 1050 840 60/70 8 90 PS* 10 PS 121210 550 120 1040 840 60/70 7 100 PS* 11 PS 13 1230 680 110 1030 84060/70 8 132 Conventional 1200 — — Ar₃ 960 Naturally — Steel 11 or moreCooled

[0156] TABLE 10 Ratios of Alloying Elements after NitrogenizingTreatment Mn/S Ti/N N/B Al/N V/N (Ti + 2Al + 4B + V)/N Present Sample 1308 1.3 15.0 3.8 1.0 10.2 Present Sample 2 308 1.2 16.4 3.5 0.9 9.3Present Sample 3 308 1.2 17.1 3.3 0.8 8.9 Comparative Sample 1 308 1.910.3 5.6 1.4 14.8 Comparative Sample 2 308 0.4 45.1 1.3 0.3 3.4 PresentSample 4 300 1.8 28.0 2.5 0.4 7.3 Present Sample 5 221 1.4 36.7 5.5 1.814.2 Present Sample 6 296 2.5 16.0 2.5 6.3 14.0 Present Sample 7 380 1.720.0 3.0 1.7 9.5 Present Sample 8 300 2.0 10.0 2.5 9.0 16.4 PresentSample 9 227 1.3 14.4 3.5 1.7 10.3 Present Sample 10 220 1.5 12.0 5.00.8 12.7 Present Sample 11 300 2.2 22.5 2.8 2.2 10.2 Present Sample 12242 2.5 16.7 4.5 2.0 13.7 Present Sample 13 368 1.4 12.0 3.6 — 8.9Conventional Steel 1 218 4.1 13.8 0.6 — 5.7 Conventional Steel 2 447 2.596.0 0.8 — 4.0 Conventional Steel 3 480 0.8 105.8 0.4 — 1.5 ConventionalSteel 4 657 4.1 4.0 0.8 8.8 15.5 Conventional Steel 5 440 6.5 4.0 1.118.5  28.1 Conventional Steel 6 980 3.2 2.6 0.4 16.1  21.6 ConventionalSteel 7 720 1.0 9.9 2.5 — 6.5 Conventional Steel 8 760 1.2 14.3 0.4 —2.2 Conventional Steel 9 655 0.8 9.1 2.1 3.9 9.2 Conventional Steel 10287 0.6 9.5 3.2 1.5 8.9 Conventional Steel 11 113 5.5 12.7 3.4 7.8 20.3

[0157] Test pieces were sampled from the hot-rolled plates manufacturedas described above. The sampling was performed at the central portion ofeach rolled plate in a thickness direction. In particular, test piecesfor a tensile test were sampled in a rolling direction, whereas testpieces for a Charpy impact test were sampled in a directionperpendicular to the rolling direction.

[0158] Using test pieces sampled as described above, characteristics ofprecipitates in each steel product (matrix), and mechanical propertiesof the steel product were measured. The results are described in Table11. Also, the microstructure and impact toughness of the heat affectedzone were measured. The results are described in Table 12. Thesemeasurements were carried out in the same fashion as in Example 1. TABLE11 Characteristics of Base Characteristics of Metal StructurePrecipitates of Volume Mechanical Properties of Matrix TiN + MnSFraction Impact Density Mean of Yield Tensile Toughness (number SizeSpacing Ferrite Thickness Strength Strength Elongation at −40° C. Samplemm²) (μm) (μm) AGS FGS (%) (mm) (MPa) (MPa) (%) (J) Present 2.3 × 10⁸0.016 0.26 17 6 82 20 454 573 35 364 Sample 1 Present 3.1 × 10⁸ 0.0170.26 15 5 84 20 395 581 36 355 Sample 2 Present 2.5 × 10⁸ 0.012 0.24 134 83 20 396 580 36 358 Sample 3 Comparative 4.3 × 10⁶ 0.154 1.4  38 27 70 20 393 584 28 212 Sample 1 Comparative 5.4 × 10⁶ 0.155 1.5  34 23  7520 392 580 29 189 Sample 2 Present 3.2 × 10⁸ 0.025 0.35 15 6 83 25 396588 35 358 Sample 4 Present 2.6 × 10⁸ 0.013 0.32 14 6 82 25 396 582 35349 Sample 5 Present 3.3 × 10⁸ 0.026 0.42 15 6 84 25 390 583 35 358Sample 6 Present 4.6 × 10⁸ 0.024 0.45 16 5  3 30 390 584 35 346 Sample 7Present 4.3 × 10⁸ 0.014 0.35 15 6 82 30 392 582 36 352 Sample 8 Present5.6 × 10⁸ 0.028 0.36 15 6 81 30 391 586 36 348 Sample 9 Present 5.2 ×10⁸ 0.021 0.35 15 8 82 30 394 586 35 358 Sample 10 Present 3.7 × 10⁸0.029 0.29 14 7 84 35 390 596 36 362 Sample 11 Present 3.2 × 10⁸ 0.0250.25 16 8 83 35 396 582 35 347 Sample 12 Present 3.2 × 10⁸ 0.024 0.34 156 87 35 387 568 36 362 Sample 13 Present 3.2 × 10⁸ 0.025 0.35 15 7 89 35388 559 35 350 Sample 14 Present 3.2 × 10⁸ 0.023 0.36 14 6 81 30 382 56238 364 Sample 15 Conventional 35 406 436 — Steel 1 Conventional 35 405441 — Steel 2 Conventional 25 629 681 — Steel 3 ConventionalPrecipitates of MgO-TiN 40 472 609 32 Steel 4 3.03 × 10⁶/mm²Conventional Precipitates of MgO-TiN 40 494 622 32 Steel 5 4.07 ×10⁶/mm² Conventional Precipitates of MgO-TiN 50 812 912 28 Steel 6 2.80× 10⁶/mm² Conventional 25 629 681 — Steel 7 Conventional 50 504 601 —Steel 8 Conventional 60 526 648 — Steel 9 Conventional 60 760 829 —Steel 10 Conventional 0.2 μm or less 11.1 × 10³ 50 401 514 18.3 Steel 11

[0159] Referring to Table 11, it can be seen that the density ofprecipitates (complex precipitates of TiN and MnS in each hot-rolledproduct manufactured in accordance with the present invention is1.0×10⁸/mm² or more, whereas the density of precipitates in eachconventional product is 4.07×10⁵/mm² or less. That is, the product ofthe present invention is formed with precipitates having a very smallgrain size while being dispersed at a considerably increased density.

[0160] The products of the present invention have a matrix structurehaving fine ferrite at a high fraction of 87% or more. TABLE 12Microstructure of Heat Affected Zone Mechanical Reproducible HeatAffected Zone with Heat Input of Properties of Impact Toughness (J) at−40° C. 100 kJ/cm Welded Zone (Maximum Heating Temp. 1,400° C.) GrainSize of Mean Δ t⁸⁰⁰⁻⁵⁰⁰ = Δ t⁸⁰⁰⁻⁵⁰⁰ = Δ t⁸⁰⁰⁻⁵⁰⁰ = Austenite in HeatVolume Grain 180 sec 120 sec 180 sec Affected Zone (μm) Fraction Size ofYield Tensile Impact Transition Impact Transition 1,200 1,300 1400 ofFerrite Ferrite Strength Strength Toughness Temp. Toughness Temp. Sample(° C.) (° C.) (° C.) (%) (μm) (kg/mm²) (kg/mm²) (J) (° C.) (J) (° C.) PS1 23 33 56 73 16 370 −74 330 −67 294 −62 PS 2 22 34 55 76 15 383 −76 353−69 301 −63 PS 3 23 32 56 74 17 365 −72 331 −67 298 −63 CS 1 54 84 18236 32 126 −43 47 −34 26 −27 CS 2 65 91 198 37 35 104 −40 35 −32 18 −26PS 4 25 37 65 75 18 353 −71 325 −68 287 −64 PS 5 26 40 57 74 16 362 −71333 −67 296 −61 PS 6 25 31 53 76 17 386 −73 353 −69 305 −62 PS 7 24 3455 74 18 367 −71 338 −67 293 −63 PS 8 27 36 53 73 14 364 −71 334 −67 294−61 PS 9 24 36 52 74 17 367 −72 335 −67 285 −62 PS 22 35 53 73 18 385−72 345 −66 294 −61 10 PS 26 34 64 74 16 358 −71 324 −68 285 −63 11 PS27 38 64 74 18 355 −71 324 −67 284 −62 12 PS 24 32 54 75 16 367 −72 336−68 285 −63 13 PS 25 31 58 72 17 365 −72 330 −68 280 −63 14 PS 24 32 5476 14 368 −72 345 −68 286 −63 15 CS* 187 −51 1 CS* 156 −48 2 CS* 148 −503 CS* 230 93 143 −48 132 (0° C.) 4 CS* 180 87 132 −45 129 (0° C.) 5 CS*250 47 153 −43  60 (0° C.) 6 CS* 141 −54 −61 7 CS* 156 −59 −48 8 CS* 145−54 −42 9 CS* 138 −57 −45 10  CS* 141 −43 219 (0° C.) 11 

[0161] Referring to Table 12, it can be seen that the size of austenitegrains under a maximum heating temperature of 1,400° C., as in the heataffected zone, is within a range of 52 to 65 μm in the case of thepresent invention, whereas the austenite grains in the conventionalproducts are very coarse to have a grain size of about 180 μm. Thus, thesteel products of the present invention have a superior effect ofsuppressing the growth of austenite grains at the heat affected zone ina welding process.

[0162] Where a welding process using a heat input of 100 kJ/cm isapplied, the steel products of the present invention have a ferritefraction of about 70% or more.

[0163] Under a high heat input welding condition in which a welding heatinput is 250 kJ/cm (the time taken for cooling from 800° C. to 500° C.is 180 seconds), the products of the present invention exhibit asuperior toughness value of about 280 J or more as a heat affected zoneimpact toughness at −40° C. while exhibiting about −60° C. as atransition temperature. That is, the products of the present inventionexhibit a superior heat affected zone impact toughness under a high heatinput welding condition.

[0164] Under the same high heat input welding condition, theconventional steel products exhibit a toughness value of about 200 J asa heat affected zone impact toughness at 0° C. while exhibiting about−60° C. as a transition temperature.

1. A welding structural steel product having fine complex precipitatesof TiN and MnS, comprising, in terms of percent by weight, 0.03 to 0.17%C, 0.01 to 0.5% Si, 1.0 to 2.5% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al,0.008 to 0.030% N, 0.0003 to 0.01% B, 0.001 to 0.2% W, at most 0.03% P,0.003 to 0.05% S, at most 0.005% O, and balance Fe and incidentalimpurities while satisfying conditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40,2.5≦Al/N≦7, 6.5≦(Ti+2Al+4B)/N≦14, and 200≦Mn/S≦400, and having amicrostructure essentially consisting of a complex structure of ferriteand pearlite having a grain size of 20 μm or less.
 2. The weldingstructural steel product according to claim 1, further comprising 0.01to 0.2% V while satisfying conditions of 0.3≦V/N≦9, and7≦(Ti+2Al+4B+V)/N≦17.
 3. The welding structural steel product accordingto claim 1, further comprising one or more selected from a groupconsisting of Ni: 0.1 to 3.0%, Cu: 0.1 to 1.5%, Nb: 0.01 to 0.1%, Mo:0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 4. The welding structural steelproduct according to claim 1, further comprising one or both of Ca:0.0005 to 0.005% and REM: 0.005 to 0.05%.
 5. The welding structuralsteel product according to claim 1, wherein complex precipitates of TiNand MnS having a grain size of 0.01 to 0.1 μm are dispersed at a densityof 1.0×10⁷ /mm² or more and a spacing of 0.5 μm or less.
 6. The weldingstructural steel product according to claim 1, wherein when a toughnessdifference between the steel product and a heat treated zone, exhibitedwhen the steel product is heated to a temperature of 1,400° C. or more,and then cooled within 60 seconds over a cooling range of from 800° C.to 500° C., is within a range of ±30 J, when a toughness differencebetween the steel product and the heat treated zone, exhibited when thesteel product is heated to a temperature of 1,400° C. or more, and thencooled within 60 to 120 seconds over a cooling range of from 800° C. to500° C., is within a range of ±40 J, and when a toughness differencebetween the steel product and the heat treated zone, exhibited when thesteel product is heated to a temperature of 1,400° C. or more, and thencooled within 120 to 180 seconds over a cooling range of from 800° C. to500° C., is within a range of 0 to 100 J.
 7. A method for manufacturinga welding structural steel product having fine complex precipitates ofTiN and MnS, comprising the steps of: preparing a steel slab containing,in terms of percent by weight, 0.03 to 0.17% C, 0.01 to 0.5% Si, 1.0 to2.5% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al, 0.008 to 0.030% N, 0.0003to 0.01% B, 0.001 to 0.2% W, at most 0.03% P, 0.003 to 0.05% S, at most0.005% O, and balance Fe and incidental impurities while satisfyingconditions of 1.2≦Ti/N≦2.5, 10≦N/B≦40, 2.5≦Al/N≦7, 6.5≦(Ti+2Al+4B)/N≦14,and 220≦Mn/S≦400; heating the steel slab at a temperature ranging from1,000° C. to 1,250° C. for 60 to 180 minutes; hot rolling the heatedsteel slab in an austenite recrystallization range at a thicknessreduction rate of 40% or more; and cooling the hot-rolled steel slab ata rate of 1° C./min to a temperature corresponding to ±10° C. from aferrite transformation finish temperature.
 8. The method according toclaim 7, wherein the slab further contains 0.01 to 0.2% V whilesatisfying conditions of 0.3≦V/N≦9, and 7≦(Ti+2Al+4B+V)/N≦17.
 9. Themethod according to claim 7, wherein the slab further contains one ormore selected from a group consisting of Ni: 0.1 to 3.0%, Cu: 0.1 to1.5%, Nb: 0.01 to 0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 10. Themethod according to claim 7, wherein the slab further contains one orboth of Ca: 0.0005 to 0.005% and REM: 0.005 to 0.05%.
 11. The methodaccording to claim 1, wherein the preparation of the slab is carried outby adding, to molten steel, a deoxidizing element having a deoxidizingeffect higher than that of Ti, thereby controlling the molten steel tohave a dissolved oxygen amount of 30 ppm or less, adding, within 10minutes, Ti to have a content of 0.005 to 0.02%, and casting theresultant slab.
 12. The method according to claim 11, wherein thedeoxidation is carried out in the order of Mn, Si, and Al.
 13. Themethod according to claim 11, wherein the molten steel is cast at aspeed of 0.9 to 1.1 m/min in accordance with a continuous castingprocess while being weak cooled at a secondary cooling zone with a waterspray amount of 0.3 to 0.35 l/kg.
 14. A method for manufacturing awelding structural steel product having fine complex precipitates of TiNand MnS, comprising the steps of: preparing a steel slab containing, interms of percent by weight, 0.03 to 0.17% C, 0.01 to 0.5% Si, 1.0 to2.5% Mn, 0.005 to 0.2% Ti, 0.0005 to 0.1% Al, at most 0.005 N, 0.0003 to0.01% B, 0.001 to 0.2% W, at most 0.03% P, 0.003 to 0.05% S, at most0.005% O, and balance Fe and incidental impurities while satisfying acondition of 220≦Mn/S≦400; heating the steel slab at a temperatureranging from 1,000° C. to 1,250° C. for 60 to 180 minutes whilenitrogenizing the steel slab to control the N content of the steel slabto be 0.008 to 0.03%, and to satisfy conditions of 1.2≦Ti/N≦2.5,10≦N/B≦40, 2.5≦Al/N≦7, and 6.5≦(Ti+2Al+4B)/N≦14; hot rolling thenitrogenized steel slab in an austenite recrystallization range at athickness reduction rate of 40% or more; and cooling the hot-rolledsteel slab at a rate of 1° C./min to a temperature corresponding to ±10°C. from a ferrite transformation finish temperature.
 15. The methodaccording to claim 14, wherein the slab further contains 0.01 to 0.2% Vwhile satisfying conditions of 0.3≦V/N≦9, and 7≦(Ti+2Al+4B+V)/N≦17. 16.The method according to claim 14, wherein the slab further contains oneor more selected from a group consisting of Ni: 0.1 to 3.0%, Nb: 0.01 to0.1%, Mo: 0.05 to 1.0%, and Cr: 0.05 to 1.0%.
 17. The method accordingto claim 14, wherein the slab further contains one or both of Ca: 0.0005to 0.005% and REM: 0.005 to 0.05%.
 18. The method according to claim 14,wherein the preparation of the slab is carried out by adding, to moltensteel, a deoxidizing element having a deoxidizing effect higher thanthat of Ti, thereby deoxidizing the molten steel to have a dissolvedoxygen amount of 30 ppm or less, adding, within 10 minutes, Ti to have acontent of 0.005 to 0.02%, and casting the resultant slab.
 19. Themethod according to claim 18, wherein the deoxidation is carried out inthe order of Mn, Si, and Al.
 20. A welded structure having a superiorheat affected zone toughness, manufactured using a welding structuralsteel product according to any one of claims 1 to 6.